Low temperature hardenable steels with excellent machinability

ABSTRACT

The present invention relates to the application of at least partially bainitic or interstitial martensitic heat treatments on steels, often tool steels or steels that can be used for tools. The first tranche of the heat treatment implying austenitization is applied so that the steel presents a low enough hardness to allow for advantageous shape modification, often trough machining. Thus a steel product is obtained which can be shaped with ease and whose hardness can be raised to a higher working hardness with a simple heat treatment at low temperature (below austenitization temperature).

CROSS REFERENCE APPLICATIONS

This application is a continuation of U.S. patent application Ser. No.14/399,289 filed Nov. 6, 2014, which is a 371 of PCT/EP2013/059471 filedMay 7, 2013, which claims the benefit of European Patent Office (EPO)12166948.5 filed May 7, 2012, the content of which are incorporated byreference.

FIELD OF THE INVENTION

The present invention relates to the application of fully and/orpartially bainitic or interstitial martensitic heat treatments oncertain steels, often tool steels or steels that can be used for tools.The first tranche of the heat treatment implying austenitization isapplied so that the steel presents a low enough hardness to allow foradvantageous shape modification, often trough machining. But thehardness can then also be raised to the working hardness with a simpleheat treatment at low temperature (below austenitization temperature).

SUMMARY

Tool steels often require a combination of different properties whichare considered opposed. A typical example can be the yield strength andtoughness. For most tool steels the best compromise of such propertiesis believed to be obtainable when performing a purely martensitic heattreatment followed by the adequate tempering, to attain the desiredhardness.

For heavy sections it is often impossible to attain pure martensiticmicrostructure through the whole cross-section, and very often it is noteven possible to attain such a microstructure at the surface. Mixedmicrostructures with bainite and martensite have a particularly lowfracture toughness which is very detrimental for several applications,like for example those where thermal fatigue is a dominant failuremechanism.

For most tool steels to attain a martensitic microstructure trough aheavy section implies the employment of very severe cooling that caneasily lead to cracking.

The conventional way to manufacture a die comprises the following steps:

-   -   Tool steel rough machining.    -   Stress relieving.    -   Finalization of the rough machining.    -   Heat treatment    -   Final machining    -   Surface treatment (Nitriding, carburizing . . . ) and/or        coating.

Dies not requiring very high wear resistance can skip the last step.When the geometry of the die is simple, often the stress-relieving stepis skipped. For some not so demanding applications, it is customary andeconomically advantageous to use pre-hardened tool steels, thus avoidingheat treatment and proceeding to final machining right away. This isespecially interesting for big dies since the cost of the heat treatmentis proportional to the weight and the distortion associated to the heattreatment and thus mandatory final machining in hard condition isproportional to the size of the die. Also often this route is chosen dueto the time saving in the execution of the project; at least one and ahalf weeks can be saved when proceeding in this way. The biggesthandicap is that the pre-hardening hardness cannot be all too high sincethen the machining would be very costly, usually hardness below 45 HRcare chosen. It is interesting to notice that the final machining takesplace at the final hardness level, where machining is usuallyconsiderably more resource consuming. Also for many applications, thoughit would be nice to benefit from the shortened implementation time andavoid costs associated to heat treatment, it is not possible to usepre-hardened tool steels because the application demands considerablyhigher bulk hardness.

With the improvement of machining capabilities in the last years, themachining of tool steels up to 40 HRc and even 45 HRc if they have somemachinability enhancement additives or a fine, but not extremely tough,microstructure is present. In fact most pre-heat treated tool steels liein the 30-40 HRc range with some special applications tool steels in the40-45 HRc range. Indeed annealed tool steels are normally quite softeroften below 250 HB, but the difference in the machinability is not sobig. As mentioned many applications require though bulk hardness above48 HRc. In cases where a bulk hardness below 45 HRc is sufficient, but ahigher surface hardness is desirable, which happens quite often,Pre-hardened tool steels are often nitrided. For many years it has beenrealized, and is one of the big advantages of tool steels, that it isdesirable to have the tool steel soft when it is machined, and hard whenit has to work. It should be as soft as possible when machining, but upto 40 HRc or even 45 HRc is acceptable, and sufficiently hard whenworking (the optimal hardness level is application dependent). For manyapplications the optimal working hardness falls in the 48-58 HRc range.Therefore often an increase of 10-20 HRc in the “hardening” process issufficient for many applications.

In most applications, hardness is not the only relevant materialproperty for the tool steel, but some other properties are as relevantor at least relevant enough to be taken into account when designing thetooling solution. Such properties can be: toughness (resilience orfracture toughness), resistance to working conditions (corrosionresistance, wear resistance, oxidation resistance at high temperatures,. . . ), thermal properties (thermal diffusivity, thermal conductivity,specific heat, heat expansion coefficient, . . . ), magnetic and/orelectric properties, temperature resistance and many others. Often theseproperties are microstructure dependent and thus will be modified duringheat treatment. So heat treatment is optimized to render the bestproperty compromise for a given application.

There are some tool steels, or better-named special alloys, which useprecipitation hardening as one of the main hardening mechanisms togetherwith solid solution and sometimes ni-martensite. On some of those toolsteels the softest possible state is the solubilized or solutionannealed state which often lies around 30-40 HRc, and the heat treatmentapplied is a low temperature precipitation often rendering a 8-20 HRchardness increase which is sufficient for many applications asexplained. This low temperature precipitation has the advantage of oftenhaving a small and controllable distortion associated. The problem ofthose special alloys that can be substitutes for tool steels, are mainlythe low wear resistance and the very high alloy manufacturing cost. Alsotheir machinability is worse than that of a tool steel at the samehardness level mainly due to the extended usage of solid solution as ahardening mechanism.

Wear in material shaping processes is, primarily, abrasive and adhesive,although sometimes other wear mechanisms, like erosive and cavitative,are also present. To counteract abrasive wear hard particles aregenerally required in tool steels, these are normally ceramic particleslike carbides, nitrides, borides or some combination of them. In thisway, the volumetric fraction, hardness and morphology of the named hardparticles will determine the material wear resistance for a givenapplication. Also, the use hardness of the tool material is of greatimportance to determine the material durability under abrasive wearconditions. The hard particles morphology determines their adherence tothe matrix and the size of the abrasive exogenous particle that can becounteracted without detaching itself from the tool material matrix. Thebest way to counteract the adhesive wear is to use FGM materials(functionally graded materials), normally in the form of ceramic coatingon the tool material. In this case, it is very important to provide agood support for the coating which usually is quite brittle. To providethe coating with a good support, the tool material must be hard and havehard particles. In this way, for some industrial applications, it isdesirable to have a tool material with high thermal diffusivity at arelatively high level of hardness and with hard particles in the form ofsecondary carbides, nitrides and/or borides and often also primary hardparticles (in the case to have to counteract big abrasive particles).

In some applications the resistance to the working environment is morefocused on corrosion or oxidation resistance than wear although bothoften co-exist. In such cases oxidation resistance at the workingtemperature or corrosion resistance against the aggressive agent aredesirable. For such applications corrosion resistance tool steels areoften employed, at different hardness levels and with different wearresistances depending on the application.

Thermal gradients are the cause of thermal shock and thermal fatigue. Inmany applications steady transmission states are not achieved due to lowexposure times or limited amounts of energy from the source that causesa temperature gradient. The magnitude of thermal gradient for toolmaterials is also a function of their thermal conductivity (inverseproportionality applies to all cases with a sufficiently small Biotnumber).

Hence, in a specific application with a specific thermal flux densityfunction, a material with a superior thermal conductivity is subject toa lower surface loading, since the resultant thermal gradient is lower.The same applies when the thermal expansion coefficient is lower and theYoung's modulus is lower.

Traditionally, in many applications where thermal fatigue is the mainfailure mechanism, as in many casting or light alloy extrusions cases,it is desirable to maximize conductivity and toughness (usually fracturetoughness and CVN).

Most forging applications use hardness in the 48-54 HRc range, plasticinjection molding is preferably executed with tools having a hardnessaround 50-54 HRc, die casting of zink alloys is often performed withtools presenting a hardness in the 47-52 HRc range, hot stamping ofcoated sheet is mostly performed with tools presenting a hardness of48-54 HRc and for uncoated sheets 54-58 HRc. For sheet drawing andcutting applications the most widely used hardness lies in the 56-66 HRcrange. For some fine cutting applications even higher hardness are usedin the 64-69 HRc.

STATE OF THE ART

Interrupted bainitic heat treatments have been used in JP1104749 (A) fora family of tool steels where special care has been taken to try toavoid the coarse precipitation of cementite, and its associatedbrittleness, trough the addition of Al. In the present invention thehardening and tempering does also imply some geometric transformation,normally trough machining, in between the complete process but toughnessis either managed at lower levels for some applications or the strategyof having a higher degree of replacement of cementite trough othercarbides is pursued. On top in the present invention solutions withconsiderably higher corrosion resistance, thermal conductivity, wearresistance, economic advantage and/or toughness are achieved.

The effect of having a lower hardness for machining and a higher one forworking and being able to go from the lower hardness to the higherhardness with a low temperature (below austenitization) heat treatmentis often used in the so called precipitation hardening steels. Thosesteels are characterized by having an austenitic, even ferritic,substitutional martensite or even low carbon interstitial martensiticmicrostructure where the precipitates nucleate and grow to the desiredsize during the heat treatment to provide the increase in hardness andmechanical strength. Many such steels exist, as an example could bementioned the maraging steels, precipitation hardening tool steels likein U.S. Pat. No. 2,715,576, JP1104749 or the well-known Daido SteelLimited NAK55 and NAK80. The differences of such steels from the steelsof the present invention is the whole conception, microstructures used,which in this case reflect mostly even in the compositional rangesemployed and temperatures employed for the heat treatments.

SUMMARY OF THE INVENTION

The authors have discovered that the problem of having a low enoughhardness during the machining and then having the desired combination ofrelevant properties for the given application comprising a higherhardness, without having to austenitize the tool steel at hightemperatures, can be solved with a steel with the features of claim 1and a method for manufacturing steel with the features of claim 21.Inventive uses and preferred embodiments follow from the other claims.

By applying a bainitic or partially bainitic heat treatment to a toolsteel presenting a large enough secondary hardness peak, and supplyingfor machining the tool steel after quenching or with one or moretempering cycles at temperatures below the temperature where the maximumhardness peak occurs, rendering a low enough hardness for the machiningcan be generated. And after the machining, or part of it, applying atleast one stress relieving, nitriding or tempering at a temperaturebelow austenitizing temperature, delivers the desired hardness.

Alternatively a martensitic heat treatment can be performed. This isadvantageous if the hardness gradient between the lowest point beforethe secondary hardness peak and the maximum secondary hardness is big.

One additional advantage of bainitic heat treatments is that they can beattained with a less abrupt quenching rate. Also for some tool steelsthey can deliver a similar microstructure trough a thicker section. Forsome tool steels with a retarded bainitic transformation it is possibleto attain a perfectly homogeneous bainitic microstructure trough anextremely heavy section.

Bainite can be very fine and deliver high hardness and toughness if thetransformation occurs at low enough temperatures. Many applicationsrequire high toughness, whether resilience or fracture toughness. Inplastic injection applications often thin walls (in terms of resistantcross-section) are subjected to high pressures. When those walls aretall a big moment is generated on the base that often has a smallradius, and thus high levels of fracture toughness are required. In hotworking applications, the steels are often subjected to severe thermalcycling, leading to cracks on corners or heat checking on the surface.To avoid the fast propagation of such cracks it is also important forthose steels to have as high as possible fracture toughness at theworking temperature. Many efforts have been placed to attain purelymartensitic structures in such applications, either through properalloying to delay bainitic transformation kinetics, or through thedevelopment of methods to increase the cooling rate but avoidingcracking. The authors have observed that what is quite detrimental fortoughness, and especially fracture toughness is the mixture ofmartensite and bainite, even for small quantities of the latter. But ifbainite is the only phase present, or at least the dominant phase, andespecially if the bainite is a fine lower bainite then very high valuesof toughness can be attained, also fracture toughness at hightemperatures. The authors have also observed that even for higher andcoarser bainite, when the alloying level is high enough and the propertempering strategy is followed, then most of the coarse cementite can bereplaced by finer carbides and good toughness values achieved especiallyat higher temperatures. As mentioned, martensitic heat treatments areoften difficult to attain for heavy sections, or they might involvealloying which is detrimental for other properties.

The inventors have realized that a very convenient way to have amaterial that can be easily shaped and yet presents a high workinghardness without the unforeseeable deformations associated to quenchingconsists on the manufacture of a steel, often a tool steel or a steelthat can be used to build tools, delivered in a condition such thatafter the delivery the bulk hardness can be raised through a heattreatment comprising temperatures below austenitization and notrequiring any particularly fast cooling. The delivery condition willcomprise an interstitial martensitic and/or partially bainitic or any ofthe above but partially tempered microstructure.

BRIEF DESCRIPTION OF THE DRAWING

FIG. 1 is a tempering graph, where hardness evolution is plotted againsttemperature.

DETAILED DESCRIPTION OF THE INVENTION

It is possible within the present invention to obtain tool steels or anysteel that has to undergo a machining process prior to its applicationin a condition where it is easy to machine and then be able to transformit to a microstructure of higher performance by applying a heattreatment that involves only temperatures below austenitizationtemperature and no requirements for a fast cooling rate, providing thena controllable, and small distortion.

Tools are often machined from pre-heated tool steels, especially bigtools where the production cost of the tool plays a big role. Since inmany cases large amounts of machining are involved it is important forthe pre-hardened tool steels to have good machinability. For thispurpose, these steels have often elements added to enhance machinabilitylike S, Ca, Bi and even Pb. Moreover they present often an homogeneousmicrostructure in the sense of size and distribution of carbides. Mostimportantly the hardness levels to which they are pre-hardened are thosewhere machining can be carried out at fast stock removing speeds.Although machining techniques do not cease to improve, and thus thehardness level for which fast stock removal is still possible continuesto increase, a good general hardness level would be <40 HRc for veryfast machinability and rarely levels of 45 HRc are exceeded. Probably 48HRc would the maximum reasonable limit. For many applications though, 40HRc (respectively 45 HRc or even 48 HRc) are not sufficient andpre-hardened steels are associated to not excessively highproductivities for many applications. For applications requiring highermechanical properties, a different route is normally employed, whichnormally implies higher costs for the manufacturing of the die, that areafterwards recovered through the higher performance (often in terms ofdurability) of the die. This route implies a rough machining in annealedstate, where the material is soft, heat treatment and final machining(mandatory to compensate the distortions occurred during heattreatment). The final machining occurs with the material already hardand thus is comparatively more difficult and costly.

Some pre-hardened tool steels are chosen to have a high enough temperingtemperature at which the hardness is fixed so that afterwardssuperficial treatments or even coatings can be applied at lowertemperatures (to avoid distortion and loss of hardness), in such a wayincreasing the tribological performance of the die. The tool steelaccording to the present invention benefits from the advantages of bothmanufacturing routes. The tool steel is provided as a pre-hardened toolsteel in terms of hardness for fast stock removal during machining andthen the material is brought to a state of superior hardness but withoutthe uncontrolled distortion of a quenching process. What is required toattain the hardness increase is a temper-like heat treatment. Sincenormally not hardness alone will be a relevant property different heattreatment combinations will be desirable for every tool steel where thepresent invention is applicable (heat treatment combination refers tothe lower hardness treatment performed before delivery, and the underaustenitization temperature treatment or treatments performedafterwards). For some of these combinations the deformation associatedto the last part of the treatment is either small or with a high enoughreproducibility to not necessarily require any dimensional correctingmachining at a high hardness level. In such cases the treatment bringingthe steel to the high performance level, or part of it might be made asa consequence of another necessary process like a nitriding, coating,stress relieving . . . . It is also possible especially for pieces withheavy machining to make coincide the treatment with a stress relievingwhile leaving some extra stock for machining in a higher hardnesscondition (to correct possible unpredictable deformations due to thefiber cutting during the machining.

Advantageously, the tool steel or steel usable for tooling, or steel ingeneral, have a secondary hardness maximum in the tempering curve with asignificantly lower hardness at a given lower tempering temperaturepoint. For the steels of the present invention, this maximum hardnessgradient between the maximum secondary hardness peak in the temperingcurve and the point of minimum hardness at lower tempering temperaturethan the tempering temperature leading to the secondary hardness peak,should be usually at least 4 HRc, often more than 7 HRc, preferably morethan 8 HRc, even more preferably at least 10 HRc. For applications wherethe end hardness is quite high, it is desirable, and can also beattained within the present invention when following the indicatedsteps, to have a hardness gradient, as above described, of at least 15HRc and preferably more than 18 HRc or even more than 20 HRc.

The present invention is especially interesting for a broad range ofapplications when the hardness can be raised with a low temperature(below austenitization) heat treatment, acting as tempering. For mostapplications a hardness above 48 HRc is desirable. For applicationsrequiring high mechanical resistance normally 50 HRc or even 52 HRcshould be attainable, for applications with high superficial pressures(like for example when wrinkling occurs in cold or hot drawingapplications) 54 HRc or even 56 HRc should be attainable. And forcutting and drawing applications often more than 60 HRc, and even morethan 62 HRc are desirable. Applications with high wear might requireeven higher hardness above 64 HRc and even above 67 HRc. These hardnesslevels can be attained within the present invention, when following theindicated steps.

The present invention is based on a combination of alloying and properlychosen microstructures. Very significant are also the heat treatmentsand how those heat treatments are applied. For many applications of thepresent invention, the preferred microstructure is predominantlybainitic, at least 50% vol %, preferably 65% vol %, more preferably 76%vol % and even more preferably more than 92% vol %, since is normallythe type of microstructure easier to attain in heavy sections and alsobecause is the microstructure normally presenting the highest secondaryhardness difference upon proper tempering.

For some applications, especially those requiring heavy sections withmaterials presenting limited hardenability in the bainitic regime, HighTemperature bainite will be preferred since it is the first bainite toform when cooling the steel after austenitization. In this document HighTemperature bainite refers to any microstructure formed at temperaturesabove the temperature corresponding to the bainite nose in the TTTdiagram but below the temperature where the ferritic/perlitictransformation ends, but it excludes lower bainite as referred in theliterature, which can occasionally form in small amounts also inisothermal treatments at temperatures above the one of the bainiticnose. For the applications requiring high easy hardenability, the hightemperature bainite should be the majority type of bainite and thus fromall bainite is preferred at least 50% vol %, preferably 65% vol %, morepreferably 75% vol % and even more preferably more than 85% vol % to beHigh Temperature Bainite. As it is well known in metallurgical terms,bainite is one of the decomposition products when austenite is notcooled under thermodinamical equilibrium. It consists of a finenon-lamellar structure of cementite and dislocation-rich ferrite platesas it is a non-diffusion process. The high concentration of dislocationsin the ferrite present in the bainite makes this ferrite harder than itwould normally be. Often high temperature bainite will be predominantlyUpper Bainite, which refers to the coarser bainite microstructure formedat the higher temperatures range within the bainite region, to be seenin the TTT temperature-time-transformation diagram, which in turn,depends on the steel composition. The inventors have found that a way toincrease the toughness of the High Temperature Bainite, including theUpper Bainite is to reduce the grain size, and thus for the presentinvention when Tough Upper Bainite is required, grain sizes of ASTM 8 ormore, preferably 10 or more and more preferably 13 or more areadvantageous. The inventors have also seen that surprisingly high valuesof toughness can be attained with High Temperature Bainite when usingmicrostructures where cementite has been suppressed, strongly reducedand/or its morphology altered to finer lamella or even more so when thecementite is globulized. For bainites including retained austenite, thesame applies for the morphology of the retained austenite phase. This iswhat is referred as Tough High Temperature Bainite in this application:small grain size high temperature bainite and/or low cementite bainiteand/or fine lamella or globular morphology high temperature bainite. Forsome applications it is clearly preferred to have most of the hightemperature bainite being tough high temperature bainite at a volumefraction of more than a 60%, preferably more than 78%, and even morepreferably more than 88% in volume percent. The inventors have foundthat specially for low % Si alloys (lower than 1%, especially lower than0.6% and even more specially lower than 0.18%), high contents ofglobular bainite provide very high resilience which is of high interestfor several applications. In this case it is desirable to have 34% ofall bainite or more to be of globular morphology, preferably 55% ormore, more preferably 72% or more and more preferably 88% or more. Insome instances it is even possible to have all bainite having a globularmorphology. When combined with small grain size as described above forthe High Temperature Bainite in general, even unexpected high values offracture toughness can be attained. For some applications having someferrite and or perlite is not too detrimental, so for most applicationsno ferrite/perlite will be desirable or at the most a 2% or eventually a5%. The applications more tolerant to ferrite/perlite can allow up to a10% or even a 18%.

In a bainitic microstructure generally the presence of martensite leadsto a decrease in fracture toughness, for applications where fracturetoughness is not so important there are no restrictions on the fractionof bainite and martensite, but the applications where fracture toughnessmatters on predominantly bainitic microstructures will prefer theabsence of martensite or at most its presence up to a 2% or possibly upto 4%. For some compositions 8% or even 17% of martensite might betolerable and yet maintaining a high fracture toughness level.

If high fracture toughness at lower temperatures is desirable, in heavycross sections, there are two possible strategies to be followed for thesteels of the present invention within the predominantly bainitic heattreatments. Either alloy the steel to assure the martensitictransformation temperature is low enough (normally lower than 400° C.,preferably lower than 340° C., more preferably lower than 290° C. andeven lower than 240° C. For extremely fine bainite, but often associatedwith very slow transformation kinetics, the transformation temperatureshould be below 220° C., preferably below 180° C. and even below 140°C., and all transformation kinetics to stable and not so desirablestructures (ferrite/perlite, upper bainite) should be slow enough (atleast 600 seconds for 10% ferrite/perlite transformation, preferablymore than 1200 seconds for 10% ferrite/perlite transformation, morepreferably more than 2200 seconds for 10% ferrite/perlite transformationand even more preferably more than 7000 seconds for 10% ferrite/perlitetransformation. Also more than 400 seconds for 20% transformation intobainite, preferably more than 800 seconds for 20% bainite, morepreferably more than 2100 seconds for 20% bainite and most preferablyeven more than 6200 seconds for 20% bainite).

Alternatively the alloying content regarding elements with higherpropensity than Fe to alloy with % C, % N and % B has to be chosen to behigh enough. Elements having an affinity for carbon higher than iron areHf, Ti, Zr, Nb, V, W, Cr, Mo as most important ones and will be referredin this document as strong carbide formers (special attention has to beapplied since this definition does not coincide with the most common onein the literature where often Cr, W and even Mo and V are often notreferred as strong carbide formers). Elements with higher carbonaffinity than Fe will form their respective carbides or a combination ofthem before the iron carbide can form, from now on referred to asalloyed carbides. Depending on the carbide itself, properties can vary.Special cases are later on and depending on the particular propertiessought, properly described. In this sense, most significant are thepresence of % Moeq, % V, % Nb, % Zr, % Ta, % Hf, to a lesser extent % Crand all other carbide formers. Often more than 4% in weight in the sumof elements with higher affinity for carbon than iron will be present,preferably more than 6.2%, more preferably more than 7.2% and even morethan 8.4%. Given the high secondary hardness peak provided by % Moeq,often more than 4.2%, preferably more than 5.2% and even more than 6.2%will be present for a preferred embodiment of the invention. In the sameway % V can be employed and often more than 0.2% is used, preferablymore than 0.6%, more preferably more than 2.4% and most preferably evenmore than 8.4%. Finally if primary carbides are not detrimental for theapplication and cost allows, very strong carbide formers (% Zr+% Ta+%Nb+% Hf) will be used in an amount exceeding 0.1%, preferably 0.3% andmost preferably even 0.6%. It is convenient that at least 30% vol % ofthe carbides, preferably 35% vol %, more preferably 40% vol % and evenmore preferably more than 45% vol % of carbides have at least 50% at %,preferably 55% at %, more preferably 60% at % and even more preferablymore than 75% at % iron of all metallic constituents of the carbides.This allows for the desired hardness increase after the application ofthe low temperature (below AC1) heat treatment process, usually carriedout at the end user's side.

Additionally any thermo-mechanical treatment leading to a refining ofthe final grain size is advantageous, especially for predominantlybainitic heat treatments because then the effect is not only theimprovement of toughness but also in the increase of hardenability. Thesame applies for treatments avoiding carbide precipitation on grainboundaries. Such a treatment can be, for example, a first step at hightemperatures above 1.020° C. to coarsen the austenite grain size (sinceit is a diffusion process the higher the temperature is, the lower isthe time required, strain can also be introduced trough mechanicaldeformation but recrystallization avoided at this point). Then the steelis cooled fast enough to avoid transformation into stablemicrostructures (ferrite/perlite, and also bainite as much as possible)and also to minimize carbide precipitation. Finally the steel is stressreleased at a temperature close to Ac1. This will promote the nucleationof very fine grains in the final heat treatment, especially if it ispredominantly bainitic. Predominantly martensitic structures can also bedesirable in the present invention if the secondary hardness peak ishigh enough to enable for a low hardness machining and afterwardssignificant rising of the hardness upon tempering. Predominantly“martensitic structures” refers to a microstructure consisting of atleast 50% vol % interstitial martensite, preferably 65% vol %interstitial martensite, more preferably 78% vol % interstitialmartensite and even more preferably more than 88% vol % interstitialmartensite. Retained austenite can also lead to a desirable hardnessincrease upon decomposition during a tempering process. Thistransformation is not the most desirable but it can be used in thepresent invention for some applications where the rather uncontrolledvolume change associated is not too critical. If little retainedaustenite is present then the effect of its decomposition is small andthus has to be necessarily supplemented by the precipitation orseparation of alloyed carbides. Alloyed carbides are those with a highamount of metallic elements which are stronger carbide builders thaniron (more than 42% at %, preferably more than 62% at % and even morepreferably more than 82% at % of the total amount of metallicconstituents of the carbide), in the sense already described. Thus whenretained austenite is present in an amount of less than 2.9%,particularly less than 2.5% and even more so less than 1.8% in vol %,then carbide formers stronger than iron have to be present in solidsolution or any other state that allows the formation of their carbidesor mixed carbides the so called in this application and often inliterature alloy carbides, without the need of re-dissolution attemperatures above Ac1. It is desirable in this case to have a 2.2% ormore, more preferably a 3% or more and more preferably a 3.8% or more inweight percent of these strong carbide formers.

If retained austenite is present in very large amounts like more than52%, particularly more than 60% and even more so when it is more than72%, then the presence of elements capable of forming alloyed carbidescan be omitted. For the in-between cases, it can be sufficient with1.2%, preferably more than 1.8% or even also more than 2.1% in weightpercent of the strong carbide formers.

Fully martensitic structures are desirable but difficult to attain forheavy sections, so normally up to a 8% or even 24% bainite can betolerated. The amounts of ferrite/perlite admissible coincide with thoseof the bainitic treatment, although the compositions will generallyvary.

There are numerous reports in the literature about the existence of verytough lower bainite under some quite restrictive conditions that lead topoor tribological performance for some applications. The inventors haveseen that this can be solved with the usage of alloyed carbides, when %C is well equilibrated as explained in more detail later. In general forthose applications it is desirable to have a 2% or more carbide formersstronger than iron, preferably a 3.2% or more, more preferably a 4.6 ormore or even a 7.6 or more. There are even fewer reports in theliterature of the existence of tough bainite structures in the hightemperature bainite regime, like for example globular or globalizedbainite, and it is always associated to low % C contents, normally inthe range of % C<0.2 in weight percent. While this structure is verydesirable for many applications in the present invention, most of thoseapplications require mechanical and tribological properties which arewith extreme difficulty attained with such low % C contents. Theinventors have seen that surprisingly in the current invention suchstructures can be attained for considerably higher % C contents. It is apeculiarity of the present invention to have simultaneously tough hightemperature bainite and more than 0.21% weight % C, preferably more than0.26%, more preferably more than 0.31%, even more preferably more than0.34%, and most preferably even more than 0.38%. The way this isachieved is by having some of the nominal % C—the theoretical total % Cof the steel—not participating in the austenite to bainitetransformation. One effective way to do so is to have some of the % Cbound to carbides right before the transformation starts and during thetransformation. This can be accomplished by not dissolving all carbidesduring the austenization, or by performing a controlled cooling so thatcarbide precipitation takes place before the bainitic transformation.This strategy can also be employed when lower % C martensite isdesirable. In this sense, it is advantageous for some applications ofthe present invention to have 5% or more of the nominal weight % C inthe form of carbides formed before the bainitic and/or martensitictransformation, preferably 8% or more, more preferably 12% or more andeven 23% or more. Given that carbon formation is not the only way toinhabilitate it during the martensitic and/or bainitic transformation,it is more clear to account for the nominal % C that participates andthus gets incorporated to the martensitic and/or bainitictransformation. This is a microstructural reference, since a detailedanalysis of the microstructure provides the % C of all phases other thanthe martensite and/or bainite, which can be subtracted from the nominal% C and finally seen what percentage it represents. So for someapplications it is desired that the martensite and/or bainite accountfor less than 88% of the nominal C % of the steel, preferably less than80%, more preferably less than 72% and even more preferably less than66% of the nominal C % of the steel. For some other applications it isdesired that the martensite and/or bainite account for less than 88% ofthe nominal C % of the steel, preferably less than 80%, more preferablyless than 72% and even more preferably less than 66% of the nominal C %of the untempered steel. In metallurgical terms, composition of steelsis normally given in terms of Ceq, which is defined as carbon upon thestructure considering not only carbon itself, or nominal carbon, butalso all elements which have a similar effect on the cubic structures ofthe steel, normally being B, N.

Both preferred microstructures are known as metastable microstructuresof non-equilibrium phases which form by means of non-diffusion processeswhich occur when cooling from the austenite phase faster than theequilibrium rate. Carbon placed in interstitial places from theface-centered cubic structure of austenite has not enough time to go outfrom the structure because of the fast cooling and most of it remains inthe structure inducing shear stresses which finally lead to the bainiteor martensite structure, depending on cooling rate and steelcomposition. Those structures are often rather brittle right afterquenching and one way to recover some ductility and/or toughness is bytempering them. In this application references are made to temperedmartensite (mostly interstitial) and tempered bainite, with thisterminology in this text referring to a martensite and/or bainite thathas undergone any type of heating after forming (during the quenchingprocess). This heating leads at first to a relaxation of the structure,followed by a migration of the carbon atoms (often the resultingmicrostructures are given particular names in the literature: Troostite,Sorbite . . . ), transformation of the retained austenite if present,precipitation of alloyed carbides and/or morphology change andredisolution of any type of carbides (cementite and alloyed carbidesincluded) amongst others. Which mechanisms actually take place and towhat extent depends on the steel composition, original microstructureand the temperature and time of the tempering cycles applied. So anyheating after quenching (formation of the martensite and/or bainite)leads to the tempered martensite and/or tempered Bainite as referred toin this application. Often during the implementation of the presentinvention a tempering (which might be a multiple one) takes place duringthe manufacturing of the steel, and another tempering (which again mightbe a multiple one) takes place during the usage of the steel tomanufacture a component or tool. Depending on the tempering temperatureused and time, as mentioned at the beginning of this paragraph,different amounts of carbon will be expelled and different mechanismswill be involved giving rise to different microstructures and oftenhaving an effect on the hardness of the steel. For this purpose, steelsare also often referred to their tempering graph, where hardnessevolution against temperature is plotted (see FIG. 1). Normal behaviorconsists of a drop of hardness on the first stages of tempering followedby a hardness increase if, amongst others, retained austenite and/orformation of alloyed carbides takes place. For the present invention,interest will be placed on the so-called maximum secondary hardnesspeak, which is the point in the tempering graph where this hardnessincrease reaches its maximum before hardness starts falling again due tocoarsening and/or redisolution of carbides and other precipitates.

The inventive method for manufacturing the steel product comprises thefollowing steps

-   -   (a) providing a steel composition having at least one of the        following components, all percentages being in weight percent:        -   % Ni<1% or        -   % Cr>4% or        -   % C>=0.33% or        -   % Mo>2.5% or        -   % Al<0.6% or        -   at least one of W, Zr, Ta, Hf, Nb is ≥0.01% or        -   at least one of S, P, Bi, Se, Te is ≥0.01%,    -   (b) Determining the critical temperature for the initiation of        the formation of austenite upon heating (Ac1) for the selected        composition.    -   (c) Providing a heat treatment to the steel comprising heating        up above Ac1 and cooling

Preferably the method is further characterized by a microstructureconsisting of at least 50% vol. % bainite. Other embodiments furthercomprise a microstructure consisting of at least a 50 vol. %interstitial martensite and retained austenite present in a 2.5-60%vol., and carbide formers stronger than iron present in a 2% weight ormore in solid solution. Further embodiments comprise a microstructureconsisting of at least a 50 vol. % interstitial martensite and retainedaustenite is present in less than a 2.5% vol., and carbide formersstronger than iron are present in a 3% weight or more in solid solution.

Other embodiments of the method of the present invention furthercomprise: determining the tempering graph for the steel with the appliedheat treatment, stress relieving or tempering the steel to a temperaturebelow the temperature of the maximum secondary hardness peak, machiningthe steel, applying a heat treatment consisting on heating to atemperature according to the tempering graph corresponding to a hardnessincrease of 4 HRc or more.

The present invention is especially well suited to obtain steels for thehot stamping tooling applications. The steels of the present inventionperform especially well when used for plastic injection tooling. Theyare also well fitted as tooling for die casting applications. Anotherfield of interest for the steels of the present document is the drawingand cutting of sheets or other abrasive components. Also forgingapplications are very interesting for the steels of the presentinvention, especially for closed die forging. Also for medical,alimentary and pharmaceutical tooling applications the steels of thepresent invention are of especial interest.

The present invention suits especially well when using steels presentinghigh thermal conductivity (thermal conductivity above 35 W/mK,preferably 38/mK, more preferably 42 W/mK, more preferably 48 W/mK andeven 52 W/mK), since their heat treatment is often complicatedespecially for dies with a large or complex geometry. In such cases theusage of the present invention can lead to very significant costsavings. According to a preferred embodiment of the invention, thesteel, especially the high thermal conductivity steel, can have thefollowing composition, all percentages being indicated in weightpercent:

% C_(eq) = 0.16-1.9 % C = 0.16-1.9 % N = 0-1.0 % B = 0-0.6 % Cr < 3.0 %Ni = 0-6 % Si = 0-1.4 % Mn = 0-3 % Al = 0-2.5 % Mo = 0-10 % W = 0-10 %Ti = 0-2 % Ta = 0-3 % Zr = 0-3 % Hf = 0-3 % V = 0-4 % Nb = 0-1.5 % Cu =0-2 % Co = 0-6,the rest consisting of iron and trace elements wherein,

% Ceq=% C+0.86*% N+1.2*% B,

characterized in that

% Mo+½·% W>2.0.

This composition as such forms an invention without the restrictions ofclaims 1 and 3.

In the meaning of this patent, trace elements refer to any element,otherwise indicated, in a quantity less than 2%. For some applications,trace elements are preferable to be less than 1.4%, more preferable lessthan 0.9% and sometimes even more preferable to be less than 0.78%.Possible elements considered to be trace elements are H, He, Xe, Be, O,F, Ne, Na, Mg, P, S, CI, Ar, K, Ca, Sc, Fe, Zn, Ga, Ge, As, Se, Br, Kr,Rb, Sr, Y, Tc, Ru, Rh, Pd, Ag, Cd, In, Sn, Sb, Te, I, Xe, Cs, Ba, La,Ce, Pr, Nd, Pm, Sm, Eu, Gd, Tb, Dy, Ho, Er, Tm, Yb, Lu, Re, Os, Ir, Pt,Au, Hg, Tl, Pb, Bi, Po, At, Rn, Fr, Ra, Ac, Th, Pa, U, Np, Pu, Am, Cm,Bk, Cf, Es, Fm, Md, No, Lr, Rf, Db, Sg, Bh, Hs, Mt alone and/or incombination. For some applications, some trace elements or even traceelements in general can be quite detrimental for a particular relevantproperty (like it can be the case sometimes for thermal conductivity andtoughness). For such applications it will be desirable to keep traceelements below a 0.4%, preferably below a 0.2%, more preferably below0.14% or even below 0.06%.

It should be clear that from all the possible compositions within therange only those are of interest where the microstructure described inthe present invention is attainable. Some smaller ranges within theabove mentioned compositional range are of special significance forcertain applications. For example when it comes to the % Ceq content itis preferably to have a minimum value of 0.22% or even 0.33%. On theother hand for very high conductivity applications it is better to keep% C below 1.5% and preferably below 0.9%. % Ceq has a strong effect inreducing the temperature at which martensitic transformation starts,thus higher values of % Ceq will be desirable for either high wearresistance applications or applications where a fine bainite isdesirable. In such cases it is desirable to have a minimum of 0.4% ofCeq often more than 0.5% and even more than 0.8%. If some other elementsthat reduce the martensite transformation temperature are present (likefor example % Ni) then the same effect can be obtained with lower % Ceq(same levels as described before). Also the % Moeq (% Mo+½·% W) levelsshould be higher for maximum thermal conductivity, normally above 3.0%often above 3.5%, preferably above 4% or even 4.5%. But high levels of %Moeq do tend to shorten the bainitic transformation time. Also ifthermal conductivity needs to be maximized is better to do so within acompositional range with lower % Cr, normally less than 2.8% preferablyless than 1.8% and even less than 0.3%. A special attention has to beplaced in elements that increase hardenability by slowing the kineticsof the austenite decomposition into ferrite/perlite. Very effective inthis sense is % Ni and somewhat less % Mn. Thus for heavy sections it isoften desirable to have a minimum % Ni content normally 1%, preferably1.5% and even 3%. If % Mn is chosen for this goal higher amounts arerequired to attain the same effect. About double as much quantity isrequired as is the case for % Ni. For applications where the steel is toattain temperatures in excess of 400° C. during service it might be veryinteresting to have % Co present which tends to increase temperingresistance amongst others and presents the odd effect of affecting thethermal diffusivity positively for high temperatures. Although for somecompositions an amount of 0.8% might suffice, normally it is desirableto have a minimum of 1.0% preferably 1.5% and for some applications even2.7%. Also for applications where wear resistance is important it isadvantageous to use strong carbide formers, then % Zr+% Hf+% Nb+% Tashould be above 0.2%, preferably 0.8% and even 1.2%. Also % V is a goodcarbide former that tends to form quite fine colonies but has a higherincidence on thermal conductivity than some of the former, but inapplications where thermal conductivity should be high but is notrequired to be extremely high and wear resistance and toughness are bothimportant, it will generally be used with a content above 0.1%,preferably 0.3% and most preferably even more than 0.55%. For very highwear resistance applications it can be used with a content higher than1.2% or even 2.2%. Other elements may be present, especially those withlittle effect on the objective of the present invention. In general itis expected to have less than 2% of other elements (elements notspecifically cited), preferably 1%, more preferably 0.45% and even 0.2%.

So, for such kind of steels, unusually high final tempering-liketemperatures (final tranche of the heat treatment to raise hardness) endup being used, often above 600° C., even when values for the hardnessover 50 HRc are chosen. In steels of the present invention it is usualto achieve a hardness of 47 HRc, sometimes more than 52 HRc, and oftenmore than 53 HRc and with the embodiments regarded as particularlyadvantageous due to their wear resistance, a hardness above 54 HRc, andoften above 56 HRc is possible with even one tempering cycle above 590°C., giving a low scattering structure characterized by a thermaldiffusivity higher than 8 mm²/s and often more than 9 mm²/s, or evenmore than 10 mm²/s, when particularly well executed even greater than 11mm²/s, even greater than 12 mm²/s and occasionally above 12.5 mm²/s. Aswell as achieving hardness greater than 46 HRc, even more than 50 HRcwith the last tempering cycle above 600° C., often above 640° C., andsometimes even above 660° C., presenting a low scattering structurecharacterized by a thermal diffusivity higher than 10 mm²/s, or eventhan 12 mm²/s, when particularly well executed then greater than 14mm²/s, even greater than 15 mm²/s and occasionally above 16 mm²/s. Thosealloys can present even higher hardness with lowering temperingtemperatures, but for most of the intended applications a high temperingresistance is very desirable. As can be seen in the examples with somevery particular embodiments with high carbon and high alloying, leadingto a high volume fraction of hard particles, a hardness above 60 HRcwith low scattering structures characterized by thermal diffusivityabove 8 mm²/s and generally more than 9 mm²/s are possible in thepresent invention.

According to a preferred embodiment of the present invention the steelscan have the following composition, all percentages being indicated inweight percent:

% C_(eq) = 0.15-3.0 % C = 0.15-3.0 % N = 0-1.6 % B = 0-2.0 % Cr > 4.0 %Ni = 0-6.0 % Si = 0-2.0 % Mn = 0-3 % Al = 0-2.5 % Mo = 0-15 % W = 0-15 %Ti = 0-2 % Ta = 0-3 % Zr = 0-3 % Hf = 0-3 % V = 0-12 % Nb = 0-3 % Cu =0-2 % Co = 0-6,the rest consisting of iron and trace elements wherein,

% Cq=% C+0.86*% N+1.2*% B.

This composition as such forms an invention without the restrictions ofclaims 1 and 3.

It should be clear that from all the possible compositions within therange only those are of interest where the microstructure described inthe present invention is attainable. Some smaller ranges within theabove mentioned compositional range are of special significance forcertain applications. For example when it comes to the % Ceq content itis preferably to have a minimum value of 0.22%, preferably 0.28% morepreferably 0.34% and when wear resistance is preferably 0.42% and evenmore preferably 0.56%. Very high levels of % Ceq are interesting due tothe low temperature at which martensite transformation starts. Suchapplications favor % Ceq maximum levels of 1.2%, preferably 1.8% andeven 2.8%. Applications where toughness is very important favor lower %Ceq contents, and thus maximum levels should remain under 0.9%preferably 0.7% and for very high toughness under 0.57%. Although anoticeable ambient resistance can be attained with 4% Cr, usually higherlevels of % Cr are recommendable, normally more than 8% or even morethan 10%. For some special attacks like those of chlorides it is highlyrecommendable to have % Mo present in the steel, normally more than 2%and even more than 3.4% offer a significant effect in this sense. Alsofor applications where wear resistance is important it is advantageousto use strong carbide formers, then % Zr+% Hf+% Nb+% Ta should be above0.2%, preferably 0.8% and even 1.2%. Also % V is good carbide formerthat tends to form quite fine colonies but has a higher incidence onthermal conductivity than some of the former, but in applications wherethermal conductivity should be high but is not required to be extremelyhigh and wear resistance and toughness are both important, it willgenerally be used with a content above 0.1%, preferably 0.54% and evenmore than 1.15%. For very high wear resistance applications it can beused with content higher than 6.2% or even 8.2%. Other elements may bepresent, especially those with little effect on the objective of thepresent invention. In general it is expected to have less than 2% ofother elements (elements not specifically cited), preferably 1%, morepreferably 0.45% and even 0.2%.

The steels described above can be particularly interesting forapplications requiring a steel with improved ambient resistance,especially when high levels of mechanical characteristics are desirableand the cost associated to heat treatment (both in terms of time andmoney) for its execution or associated distortions, are significant.

According to another preferred embodiment of the present invention thesteels can have the following composition, all percentages beingindicated in weight percent:

% C_(eq) = 0.15-2.0 % C = 0.15-0.9 % N = 0-0.6 % B = 0-0.6 % Cr > 11.0 %Ni = 0-12 % Si = 0-2.4 % Mn = 0-3 % Al = 0-2.5 % Mo = 0-10 % W = 0-10 %Ti = 0-2 % Ta = 0-3 % Zr = 0-3 % Hf = 0-3 % V = 0-12 % Nb = 0-3 % Cu =0-2 % Co = 0-12,the rest consisting of iron and trace elements wherein,

% C _(eq)=% C+0.86*% N+1.2*% B.

This composition as such forms an invention without the restrictions ofclaims 1 and 3.

It should be clear that from all the possible compositions within therange only those are of interest where the microstructure described inthe present invention is attainable. Some smaller ranges within theabove mentioned compositional range are of special significance forcertain applications. For example when it comes to the % Ceq content itis preferably to have a minimum value of 0.22%, preferably 0.38% morepreferably 0.54% and when wear resistance is important preferably 0.82%,more preferably 1.06% and even more than 1.44%. Very high levels of %Ceq are interesting due to the low temperature at which martensitetransformation starts, such applications favor % Ceq maximum levels of0.8%, preferably 1.4% and even 1.8%. Applications where toughness isvery important favor lower % Ceq contents, and thus maximum levelsshould remain under 0.9% preferably 0.7% and for very high toughnessunder 0.57%. Although corrosion resistance for martensiticmicrostructure can be attained with 11% Cr, usually higher levels of %Cr are recommendable, normally more than 12% or even more than 16%. Forsome special attacks like those of chlorides and to enhance hardnessgradient at the secondary hardness peak it is highly recommendable tohave % Moeq present in the steel, often more than 0.4%, preferably morethan 1.2% and even more than 2.2% offer a significant effect in thissense. Also for applications where wear resistance or thermalconductivity are important it is advantageous to use strong carbideformers, then % Zr+% Hf+% Nb+% Ta should be above 0.1%, preferably 0.3%and even 1.2%. Also % V is good carbide former that tends to form quitefine colonies but has a higher incidence on thermal conductivity thansome of the former, but in applications where thermal conductivityshould be high but is not required to be extremely high and wearresistance and toughness are both important, it will generally be usedwith a content above 0.1%, preferably 0.24% and even more than 1.15%.For very high wear resistance applications it can be used with contenthigher than 4.2% or even 8.2%. Other elements may be present, especiallythose with little effect on the objective of the present invention. Ingeneral it is expected to have less than 2% of other elements (elementsnot specifically cited), preferably 1%, more preferably 0.45% and even0.2%.

The steels described above can be particularly interesting forapplications requiring a steel with corrosion or oxidation resistance,especially when high levels of mechanical characteristics are desirableand the cost associated to heat treatment (both in terms of time andmoney) for its execution or associated distortions, are significant.

According to another embodiment of the present invention the steels canhave the following composition, all percentages being indicated inweight percent:

% C_(eq) = 0.5-3.0 % C = 0.5-3.0 % N = 0-2.2 % B = 0-2.0 % Cr = 0.0-14 %Ni = 0-6.0 % Si = 0-2.0 % Mn = 0-3 % Al = 0-2.5 % Mo = 0-15 % W = 0-15 %Ti = 0-4 % Ta = 0-4 % Zr = 0-12 % Hf = 0-4 % V = 0-12 % Nb = 0-4 % Cu =0-2 % Co = 0-6,the rest consisting of iron and trace elements wherein,

% C _(eq)=% C+0.86*% N+1.2*% B.

This composition as such forms an invention without the restrictions ofclaims 1 and 3.

It should be clear that from all the possible compositions within therange only those are of interest where the microstructure described inthe present invention is attainable. Some smaller ranges within theabove mentioned compositional range are of special significance forcertain applications. For example when it comes to the % Ceq content itis preferably to have a minimum value of 0.62%, preferably 0.83% morepreferably 1.04% and when extreme wear resistance is importantpreferably 1.22%, more preferably 1.46% and even more than 1.64%. Veryhigh levels of % Ceq are interesting due to the low temperature at whichmartensite transformation starts, such applications favor % Ceq maximumlevels of 1.8%, preferably 2.4% and even 2.8%. % Cr has two ranges ofparticular interest: 3.2%-5.5% and 5.7%-9.4%. To enhance hardnessgradient at the secondary hardness peak it is highly recommendable tohave % Moeq present in the steel, often more than 2.4%, preferably morethan 4.2% and even more than 10.2% offer a significant effect in thissense. Also for applications where wear resistance or thermalconductivity are important it is advantageous to use strong carbideformers, then % Zr+% Hf+% Nb+% Ta should be above 0.1%, preferably 1.3%and even 3.2%. Also % V is good carbide former that tends to form quitefine colonies of very hard carbides, thus when wear resistance andtoughness are both important, it will generally be used with a contentabove 1.2%, preferably 2.24% and even more than 3.15%. For very highwear resistance applications it can be used with content higher than6.2% or even 10.2%. Other elements may be present, especially those withlittle effect on the objective of the present invention. In general itis expected to have less than 2% of other elements (elements notspecifically cited), preferably 1%, more preferably 0.45% and even 0.2%.It is important for the achievement of the wear resistance to have thepresence of carbide formers stronger than iron, specially the more costeffective are more often used in a more extensive way, in particulargenerally it will be % Cr+% W+% Mo+% V+% Nb+% Zr should be above 4.0%,preferably 6.2%, more preferably 8.3% and even 10.3%.

The steels described above can be particularly interesting forapplications requiring a steel with very high wear resistance,especially when high levels of hardness are desirable and the costassociated to heat treatment (both in terms of time and money) for itsexecution or associated distortions, are significant.

According to another preferred embodiment of the present invention thesteel can have the following composition, all percentages beingindicated in weight percent:

% C_(eq) = 0.2-0.9 % C = 0.2-0.9 % N = 0-0.6 % B = 0-0.6 % Cr = 0.0-4.0% Ni = 0-6.0 % Si = 0.2-2.8 % Mn = 0.2-3 % Al = 0-2.5 % Mo = 0-6 % W =0-8 % Ti = 0-2 % Ta = 0-2 % Zr = 0-2 % Hf = 0-2 % V = 0-4 % Nb = 0-2 %Cu = 0-2 % Co = 0-6,the rest consisting of iron and trace elements wherein,

% C _(eq)=% C+0.86*% N+1.2*% B,

characterized in that

% Si+% Mn+% Ni+% Cr>2.0, or

% Mo>1.2, or

% B>2 ppm

This composition as such forms an invention without the restrictions ofclaims 1 and 3.

It should be clear that from all the possible compositions within therange only those are of interest where the microstructure described inthe present invention is attainable. Some smaller ranges within theabove mentioned compositional range are of special significance forcertain applications. For example when it comes to the % Ceq content itis preferably to have a minimum value of 0.22%, preferably 0.28%, morepreferably 3.2% and even 3.6%. Very high levels of % Ceq are interestingdue to the low temperature at which martensite transformation starts,such applications favor % Ceq maximum levels of 0.6%, preferably 0.8%and even 0.9%. % Cr has two ranges of particular interest: 0.6%-1.8% and2.2%-3.4%. Particular embodiments also prefer % Cr to be 2%. To enhancehardness gradient at the secondary hardness peak it is highlyrecommendable to have % Moeq present in the steel, often more than 0.4%,preferably more than 1.2%, more preferably more than 1.6% and even morethan 2.2% offer a significant effect in this sense. In this particularapplication of the invention the elements that mostly remain in solidsolution, the most representative being % Mn, % Si and % Ni are verycritical. It is desirable to have the sum of all elements whichprimarily remain in solid solution exceed 0.8%, preferably exceed 1.2%,more preferably 1.8% and even 2.6%. As can be seen both % Mn and % Sineed to be present. % Mn is often present in an amount exceeding 0.4%,preferably 0.6% and even 1.2%. For particular applications, Mn isinteresting to be even 1.5%. The case of % Si is even more criticalsince when present in significant amounts it strongly contributes to theretarding of cementite coarsening. Therefore % Si will often be presentin amounts exceeding 0.4%, preferably 0.6% and even 0.8%. When theeffect on cementite is pursuit then the contents are even bigger, oftenexceeding 1.2%, preferably 1.5% and even 1.65%. Also for applicationswhere wear resistance or thermal conductivity are important it isadvantageous to use strong carbide formers, then % Zr+% Hf+% Nb+% Tashould be above 0.1%, preferably 1.3% and even 2.2%. Also % V is goodcarbide former that tends to form quite fine colonies of very hardcarbides, thus when wear resistance and toughness are both important, itwill generally be used with a content above 0.2%, preferably 0.4% andeven more than 0.8%. For very high wear resistance applications it canbe used with content higher than 1.2% or even 2.2%. Other elements maybe present, especially those with little effect on the objective of thepresent invention. In general it is expected to have less than 2% ofother elements (elements not specifically cited), preferably 1%, morepreferably 0.45% and even 0.2%. As can be seen the critical elements forattaining the mechanical properties desired for such applications needto be present and thus it has to be % Si+% Mn+% Ni+% Cr greater than2.0%, preferably greater than 2.2%, more preferably greater than 2.6%and even greater than 3.2%. For some applications it is interesting toreplace % Cr for % Mo, due to the higher effect on the secondaryhardness peak and the improved thermal conductivity potential it impairsthe steel, and then the same limits apply. Alternatively to % Si+% Mn+%Ni+% Mo>2.0% . . . the presence of % Mo can be dealt alone when presentin an amount exceeding 1.2%, preferably exceeding 1.6%, and evenexceeding 2.2%. For the applications where cost is important it isspecially advantageous to have the expression % Si+% Mn+% Ni+% Crreplaced by % Si+% Mn and then the same preferential limits can apply,but in presence of other alloying elements, also lower limits can beused like % Si+% Mn>1.1%, preferably 1.4% or even 1.8%. For someapplications, % Ni is desirable to be at least 1%. For this kind ofsteels tough bainite treatments at temperatures close to martensitestart of transformation (Ms) are very interesting (often 70% or more,preferably 70% and more, or even 82% or more of the transformation ofaustenite should take place below 520° C., preferably 440° C., morepreferably 410° C. or even 380° C., but not below 50° C. belowmartensite start of transformation [Ms]). To lower the hardness formachining one or several long tempering cycles around cementiteseparation and cementite coalescence but below Chromium carbideprecipitation (alternatively Molybdenum carbide) can be used. The actualtemperature is composition dependent but often between 380 and 460° C.

The steels described above can be also applied for the manufacturing ofbig plastic injection tools particularly interesting for applicationsrequiring very low cost steel with high mechanical resistance andtoughness. This particular application of the present invention is alsointeresting for other applications requiring inexpensive steels withhigh toughness and considerable yield strength. It is particularlyadvantageous when the steel requires a harder surface for theapplication and the nitriding or coating step is made coincide with thehardening step.

A very interesting aspect of the present invention, leading tosignificant cost reductions, is given when the amount of machiningrequired in hard state can be minimized or even eliminated. This is sobecause the machining at high hardness is costly. The present inventionallows to do so, given the small amount of deformation associated tosome of the below austenitization hardening low temperature heattreatments. Most importantly the deformation is highly reproducible andisotropic for which reason it can be taken into account and compensatedfor during the machining in softer condition. The composition and heattreatment strategy has to be well chosen for the deformation during thelast tranche of the heat treatment to be small enough to avoid machiningin hard state, which allows making coincide the sub-austenitizationtemperature hardening heat treatment to coincide with the nitriding orother superficial treatment. As an illustrative example, for many of thesteels of the present invention when % Cr and % Si are low and % Moeq israther high, and when a bainitic treatment is chosen, normally thematerial will shrink for low tempering temperatures, expand close fortemperatures close to the maximum secondary hardness peak, and shrinkagain for higher temperatures, thus it is possible if the material isnot tempered or just tempered at very low temperatures, to find atemperature above the temperature delivering maximum secondary hardness,which renders almost no net deformation in the last tranche of the heattreatment (compensation of shrinkage with expansion). Thus it is aspecial execution of the present invention steels that can be deliveredwith a low enough hardness for massive machining after quenching (withor without tempering) which can suffer very slight, reproducible andisotropic deformation when the final hardness rising part of the heattreatment is applied. Thus the steel will then be characterized by anattainable deformation, in the last sub-austenitization temperaturehardening tranche of the heat treatment, smaller than 0.2% preferablysmaller than 0.1%, more preferably smaller than 0.05% and even smallerthan 0.01%. Also the difference in the deformation in two differentdirections, isotropy of the deformation, can be made to be higher than a60%, preferably higher than a 72%, often higher than 86% and even higherthan a 98%. When it comes to reproducibility, it is possible with anespecial execution of the present invention to attain reproducibility ofthe deformation in the last tranche of the hardening process above a60%, preferably above a 78%, often above a 86% and even above a 96%.(Reproducibility measured as the percentage difference of thedeformation occurred in one same orientation with two selected identicaltreatments).

Indeed one main aspect for many of the steels of the present inventionis the possibility of easily machining, even in big amounts, in a statethat does not require austenitization afterwards to attain the desiredworking hardness, and this in steels that are not precipitationhardening. Therefore it is important to have a low hardness after thefirst tranche of the treatment involving austenitization. Normally 48HRc still allow for quite fast turning, but if form milling is involvedthe hardness should not exceed 45 HRc and preferably 44 HRc and even beless than 42 HRc. If some more complex operations like honing or screwtapping have to be carried away then it is desirable that the attainablehardness can be even lower than 40 HRc, preferably 38 HRc or even lowerthan 36 HRc.

The temperatures involved in the last tranche of the heat treatment,which are always below austenitization temperature, play a significantrole for some applications. For instance, in some applications it isdesirable to have such temperature as high as possible, since thoseapplications benefit either from the tempering resistance or the higherstability associated to a high temperature tempering. Thus for thoseapplications it is desirable to have the ability to attain the workinghardness even if temperatures above 600° C., preferably 620° C., morepreferably 640° C. and even 660° C. are involved. On the other hand someapplications benefit from having the temperature for the last tranchehardening cycle at the common temperatures employed for superficial heattreatments, and especially when an acceptably low deformation or highenough deformation stability occurs with this treatment. Suchtemperatures are for example 480° C., 500° C. to 540° C. and 560° C.

One way for the steels of the present invention to be able to increasetheir hardness through a low temperature tempering like thermaltreatment, is by assuring that the right type of carbides are present atthe moment of delivery of the steel, so that it is desirable that atleast 30% vol % of all the carbides, preferably 35% vol % or more, morepreferably 42% vol % and even more preferably more than 58% vol % ofcarbides have at least 50% at %, preferably 55% at %, more preferably62% at % and even more preferably more than 73% at % iron of allmetallic constituents of the carbides. Another possible way is byassuring that at the moment of delivery the steel microstructurepresents less than 70% of the alloyed carbides, preferably less than65%, more preferably less than 58% and even less than 42% of thementioned alloyed carbides that can be attained (maximum vol % possible)with the chosen composition according to simulation for phase equilibriasoftware packages, like for example Themo-Calc or MTDATA.

The increase in hardness in the last tranche of the heat treatment ismainly attained trough the precipitation of alloy carbides, but can alsobe a consequence of the transformation of retained austenite. For manycompositions in the present invention, a separation of cementite frommartensite occurs at temperatures around 450° C. leading to a decreasein hardness often used in the present invention to provide the lowhardness machining delivery condition. This point of lowest hardness inthe tempering graph can be as low as 300° C. and as high as 540° C. Whentempering at higher temperatures in the final tranche of the heattreatment for all possible microstructures in the present inventiondissolution of the cementite and the carbon that goes into solidsolution can contribute to the separation or further precipitation ofalloyed carbides, that is carbides containing carbide forming elements.(Cr, Mo, W, V, Nb, Zr, Ta, Hf . . . ), often mixed carbides containingthose elements and others like for example iron. Those carbides oftenprecipitate as M7C3, M4C3, MC, M6C, M2C. The temperature at which thishappens is often above 400° C., preferably 450° C., more preferably 480°C. and even 540° C. Another mechanism that is profited from with somecompositions of the present invention to contribute to the hardnessincrease is the decomposition of retained austenite.

Available carbon, i.e. carbon which is not combined with any otherelement in the form of carbides and which can be found in solid solutionor not, as well as the nature of the alloyed carbides will have aneffect on the amount of hardness increase once the proper tempering isapplied.

It is clear that the present invention is especially advantageous whenabundant machining has to be undergone by the steel, and yet high bulkworking hardness is desirable. In fact the present invention isparticularly advantageous if more than a 10% of the original weight ofthe steel block has to be removed to attain the final geometry, moreadvantageous when more than 26% has to be removed, and even moreadvantageous when more than 54% has to be removed. Most machining willnormally take place between the first tranche of the heat treatmentinvolving austenitization and eventual one or more tempering-like cyclesand the final tranche of the heat treatment. In fact often at least a32% of the total machining will occur in this state, often more than 54%of the total machining, even more than 82% of the total machining whennot the 100%. In some instances it might be advantageous to perform somemachining before the part of the heat treatment involvingaustenitization, like for example long holes or any other kind ofmachining especially when it is difficult. And as mentioned beforemachining in the hard state does happen quite often, but normally insmall amounts given its higher cost.

To attain the high levels of hardness and wear resistance sometimesdesirable in the present invention, considerably high levels of thevolume fraction of hard particles have to be used. The volume fractionof hard particles (carbides, nitrides, borides and mixtures thereof) isoften above a 3%, preferably above 4.2%, more preferably above a 5.5%,and for some high wear applications, even above a 8%. Size of primaryhard particles is very important to have an effective wear resistanceand yet not excessively small toughness. The inventors have observedthat for a given volume fraction of hard particles the overallresilience of the material diminishes as the size of the hard particlesincreases, as would be expected. More surprisingly it has also beenobserved that when the size of hard particles is increased, the overallfracture toughness increases if the fracture toughness of the particlesthemselves is maintained. When it comes to abrasive wear resistance ithas been observed the existence of a critical hard particle size, belowwhich the hard particle is not effective against the abrasive agent.This critical size depends on the size of the abrasive agent and thenormal pressure. For some applications where the abrasive particles areof small size (normally below 20 microns), it can be desirable to haveprimary hard particles smaller than 10 microns or even smaller than 6microns, but in any case with an average size not smaller than 1 micron.For applications where big abrasive particles cause the wear, bigprimary hard particles will be desirable. Therefore, for someapplications it is desirable to have some primary hard particles biggerthan 12 microns, often greater than 20 microns and for some particularapplications even greater than 42 microns.

For applications where mechanical strength more than wear resistance areimportant, and it is desirable to attain such mechanical strengthwithout compromising too much toughness, the volume fraction of smallsecondary hard particles is of great importance. The term “smallsecondary hard particles” as used in the application are those with amaximum equivalent diameter (diameter of a circle with equivalentsurface as the cross section with maximum surface on the hard particle)below 7.5 nm. It is desirable to have a volume fraction of smallsecondary hard particles for such applications above 0.5%. It isbelieved that a saturation of mechanical properties for hot workapplications occurs at around 0.6%, but it has been observed by theinventors that for some applications requiring high plastic deformationresistance at somewhat lower temperatures it is advantageous to havehigher amounts than 0.6%, often more than 0.8% and even more than 0.94%.Since the morphology (including size) and volume fraction of secondarycarbides change with heat treatment, the values presented here describeattainable values with proper heat treatment.

In view of the preceding paragraphs, an effort can be made to try togroup all possible compositions of steels where the present invention isof especial interest. Of course, of all the possible compositions withinthe range only those where the microstructure described in the presentinvention is attainable are of interest. The result is that the steelwould have the following compositional restrictions:

% Ni<1% or % Cr>4% or % C≥0.33% or % Mo>2.5% or % Al<0.6% or

at least one of W, Zr, Ta, Hf, Nb, La, Ac is ≥0.01% orat least one of S, P, Bi, Se, Te is ≥0.01%

While for some steels of the present invention large quantities of % Niare desirable, for others the content has to be low enough for thepresent invention to work, in combination with the other alternativecompositional restrictions % Ni<1% is a valid limit, one would havepreferably % Ni<0.8 or even % Ni<0.2. Also for % Cr it has beenmentioned that the high thermal conductivity steels will have low % Crcontents, often below 3% and even below 0.1%, but their compositions getcovered by other alternatives in this composition, like % Mo>2.5% or %Al<0.6%, also for the ones presenting high wear resistance % C≥0.33%.But for ambient resistant steels it has to be % Cr>4%. In fact in thisglobal compositional restriction it is also preferably to have % Cr>5.3%and even % Cr>7.2%. It is also preferably to have % Mo>3.2% and evenbetter to have a restriction involving % Moeq instead of % Mo like %Moeq>2.8% or preferably % Moeq>3.4 or even % Moeq>4.2%. Anotherinteresting case is that of % Al, where it would be preferably to have %Al<0.4 or even % Al<0.16, and it would also be interesting to combinewith % Si since both are aiming at a similar goal, namely the reductionof the negative influence of Fe3C morphology on toughness. In thisrespect one could have the additional restriction with the % Alrestriction of % Si<0.8, preferably % Si<0.4 and even % Si<0.2. In thecase of carbon, it would be preferably to have % C>0.36 or even %C>0.42. It could also be possible, even convenient to make therestriction in terms of carbon equivalent instead. So one would have %Ceq≥0.33, preferably % Ceq≥0.36 or even % Ceq>0.46. In the case of theselected strong carbide formers (W, Zr, Ta, Hf, Nb, La, Ac) one wouldhave preferably more than 0.08% or even more than 0.16%. At last thecase of vanadium should be mentioned, since this element should inprinciple add two additional disjunctive restrictions, one to limit itspresence to care for high thermal conductivity steels without high wearresistance where it would be % V<1, preferably % V<0.4 and even % V<0.2.And even more important, for applications requiring high wear resistancewe should have % V>0.3, preferably % V>1.2 or even % V>3.2.

To increase machinability, S, As, Te, Bi or even Pb, Ca, Cu, Se, Sb orothers can be used, with a maximum content of 1%, with the exception ofCu that can even have a maximum content of 2%. The most commonsubstance, sulfur, has, in comparison, a light negative effect on thematrix thermal conductivity in the normally used levels to increasemachinability. However, its presence must be balanced with Mn, in anattempt to have everything in the form of spherical manganesebisulphide, less detrimental for toughness, as well as the leastpossible amount of the remaining two elements in solid solution in casethat thermal conductivity needs to be maximized. Other elements may bepresent, especially those with little effect on the objective of thepresent invention. In general it is expected to have less than 2% ofother elements (elements not specifically cited), preferably less than1%, and most preferably less than 0.45% and even less than 0.2%.

The steel of the present invention can be manufactured with anymetallurgical process, among which the most common are sand casting,lost wax casting, continuous casting, melting in electric furnace,vacuum induction melting. Powder metallurgy processes can also be usedalong with any type of atomization and eventually subsequent compactingas the HIP, CIP, cold or hot pressing, sintering (with or without aliquid phase and regardless of the way the sintering process takesplace, whether simultaneously in the whole material, layer by layer orlocalized), laser cusing, spray forming, thermal spray or heat coating,cold spray to name a few of them. The alloy can be directly obtainedwith the desired shape or can be improved by other metallurgicalprocesses. Any refining metallurgical process can be applied, like VD,ESR, AOD, VAR. Forging or rolling are frequently used to increasetoughness, even three-dimensional forging of blocks. Tool steel of thepresent invention can be obtained in any shape, for example in the formof bar, wire or powder (amongst others to be used as solder or weldingalloy). Also laser, plasma or electron beam welding can be conductedusing powder or wire made of steel of the present invention. The steelof the present invention could also be used with a thermal sprayingtechnique to apply in parts of the surface of another material.Obviously the steel of the present invention can be used as part of acomposite material, for example when embedded as a separate phase, orobtained as one of the phases in a multiphase material. Also when usedas a matrix in which other phases or particles are embedded whatever themethod of conducting the mixture (for instance, mechanical mixing,attrition, projection with two or more hoppers of different materials .. . ). The steels of the present invention can also be a part of afunctionally graded material, in this sense any protective layer orlocalized treatments can be used. The most typical ones being layers orsurface treatments:

-   -   To improve tribological performance: Superficial hardening        (laser, induction . . . ), superficial treatment (nitriding,        carburizing, borurizing, sulfidizing, any mixtures of the        previous . . . ), coatings (CVD, PVD, fluidized bed, thermal        projection, cold spray, cladding . . . ).    -   To increase corrosion resistance: hard chromium, palladium,        chemical Nickel treatment, sol gel with corrosion resistant        resins, in fact any electrolytic or non-electrolytic treatment        providing corrosion or oxidation protection.    -   Any other functional layer also when the function is appearance.

Tool steel of the present invention can also be used for themanufacturing of parts requiring a high working hardness (for exampledue to high mechanical loading or wear) which require some kind of shapetransformation from the original steel format. As an example: Dies forforging (open or closed die), extrusion, rolling. The present inventionis especially indicated for the manufacture of dies for the hot stampingor hot pressing f sheets. Dies for plastic forming of thermoplastics andthermosets in all of its forms. Also dies for forming or cutting.

EXAMPLES

Some examples indicate the way in which the steel composition of theinvention can be specified with higher precision for different hotworking applications:

Example 1

High Thermal conductivity steels (over 42 W/mK and over 8.5 mm2/s andreaching 57 W/mK and 13.5 mm2/s at 50 HRc, the thermal conductivity anddiffusivity increase for lower hardnesses at least until 40 HRc for allsteels of the present example), delivered at a hardness of 45 HRc orless and then raising the hardness to above 48 HRc after a great part ofthe machining has taken place.

For this purpose in the context of the present invention the followingcompositional range can be used:

Cq: 0.3-0.6; Cr<3.0% (preferably Cr<0.1%)

V: 0-0.9%

Si: <0.15% (preferably % Si<0.1, but with an acceptable level of oxideinclusions)

Mn: <1.0%; Mo_(eq): 2.0-8.0

where Mo_(eq)=% Mo+½% W and

C_(eq)=% C+0.86*% N+1.2*% B

The rest of the elements should be kept as low as possible and, in anycase, always be below 0.45%, with the exception of carbide formersstronger than tungsten (% Ta, % Zr, % Hf . . . ), and some solidsolution strengtheners like % Ni, % Co and eventually % Cu.

All values are given in weight percentage.

The following examples show properties that can be obtained:

Delivery Max usage Hardness hardness % C % Mo % W % V % Cr % Si % MnOther HRc HRc 0.40 3.6 1.4 0.3 <0.01 <0.05 <0.01 — 39* 56 0.32 3.36 1.910.22 <0.01 <0.05 0.4 Hf, Zr, Nb, B 41* 53 0.33 3.8 1.22 0.4 <0.01 <0.05<0.01 Hf, Zr, Nb 40* 53 0.36 3.66 1.26 0.02 <0.01 <0.05 <0.01 Zr = 0.5 37** 52 0.31 3.36 1.52 0.45 <0.01 <0.05 <0.01 Hf, Zr, Nb, Co 40* 540.36 3.75 1.91 0.44 1.12 0.1 0.47 Hf, Zr, Nb, Co 40* 55 0.32 3.36 1.11<0.01 <0.01 <0.05 <0.01 Hf, Zr, 38* 51 0.60 3.6 1.2 0.62 <0.01 0.14 0.54— 44* 58 0.72 3.75 2.0 0.54 <0.01 <0.05 <0.01 Hf, Zr, Ni, Co, B 45* 520.34 1.6 4.5 0.1 <0.01 <0.05 <0.01 Ni 2.6  38** 52 0.31 3.2 0.8 <0.01<0.01 <0.05 <0.01 Ni 0.8  37** 50 0.31 3.2 0.8 <0.01 <0.01 <0.05 <0.01Ni 0.8  47*** 52 *Delivery takes place with a mixed bainite/martensitemicrostructure where at least one tempering below 550° C. has beenapplied. **Delivery takes place with a mostly bainitic microstructurefor heavy sections and either no tempering or one or more temperingcycles under 580° C. have been applied. ***Delivery takes place with amartensitic microstructure where either no tempering or one or moretempering cycles under 580° C. have been applied.

Other Examples

Delivery Max usage Hardness Hardness % C % Mo % W % V % Cr % Si % MnOther HRc HRc 0.17 3.3 1.1 0.10 <0.01 0.2 0.36 Hf, Zr, Co 39* 50 0.652.0 <0.01 <0.01 17 0.4 0.3  44*** 51 1.23 3.8 11.2 3.4 2.01 <0.05 0.21Co  47** 62 0.98 2.66 1.26 2.02 8.01 1.05 0.17  47** 58 0.45 3.39 1.540.85 4.21 0.25 0.41 40* 51 0.61 3.34 1.65 0.52 5.08 0.32 0.32 Hf, Zr, Nb44* 57 *Delivery takes place with a mixed bainite/martensitemicrostructure where at least one tempering below 550° C. has beenapplied. **Delivery takes place with a mostly bainitic microstructurefor heavy sections and either no tempering or one or more temperingcycles under 580° C. have been applied. ***Delivery takes place with amartensitic microstructure with some perlite isles where either notempering or one or more tempering cycles under 580° C. have beenapplied.

Other Examples

Delivery Max usage Hardness Hardness % C % Mo % W % V % Cr % Si % MnOther * HRc HRc 0.29 3.36 0.1 0.002 0.019 0.04 0.022 — 40 51 0.28 3.590.6 0.003 0.02 0.04 0.025 —   40.5 53 0.28 3.70 1.19 <0.005 0.01 0.040.02 — 38 49.5 0.39 3.71 1.2 0.6 0.01 0.05 0.02 Ni 0.84, Hf, 42 53.5 Nb,Zr 0.41 3.63 1.63 0.81 0.01 0.04 0.02 Co 3.00   42.5 57 0.4 1.15 0.020.87 8.2 0.11 0.14 Ni, Al, Co 43 56 0.27 3.40 1.08 <0.005 0.01 0.05 0.02Hf 42 54 0.29 3.70 1.01 0.005 0.01 0.05 0.019 — 42 53 0.33 3.39 1.110.43 0.01 0.05 0.24 Nb 42 54 0.32 3.36 1.15 0.44 0.01 0.05 0.12 Ni 2.04338HB 53 0.29 3.62 1.18 0.004 0.01 0.05 0.02 — 40 53 0.33 3.58 1.27<0.005 0.01 0.05 0.14 Ni 3.09 41 53 0.41 3.58 1.16 0.65 0.01 0.07 0.14Nb 43 54 0.33 3.64 1.1 0.46 0.01 0.05 0.26 Nb 41 55 0.33 3.7 1.36 0.430.01 0.05 0.26 Nb, Zr 42/40 54/53.5 0.21 3.2 1.04 0.3 0.01 0.04 0.21 —42 50 0.31 3.70 2.3 <0.005 0.01 0.02 0.02 Ni 1.86 41 50 0.37 3.90 2.0<0.005 0.01 0.02 0.11 Ni 2.05 39 48.5 0.44 3.64 1.97 0.7 0.01 0.05 0.02Co 3.00 45 56 0.43 3.73 1.8 0.69 0.01 0.05 0.02 Co 3.00 44 57 0.32 3.101.68 <0.005 0.01 0.04 0.09 Ni 2.96 38 52 0.29 3.60 1.09 <0.005 0.01 0.030.015 Hf, B, Zr 42 47 0.39 3.57 1.35 0.44 <0.01 <0.01 <0.01 Hf, Zr, Nb43 53 0.32 3.1 1.7 0.030 0.1 0.1 0.17 Ni 0.017 40 50 0.356 3.900 1.4000.484 <0.01 <0.05 0.058 Ni 0.470 43 51 0.353 3.810 1.410 0.461 <0.01<0.05 0.061 Ni 0.481 137HB 53.5 0.326 3.680 1.490 0.440 0.0108 <0.050.055 Ni 0.488 40 57.5 0.464 3.890 1.670 0.452 <0.01 <0.05 0.055 Ni0.516 382HB 54.5 0.299 3.770 1.310 0.452 <0.01 <0.05 0.051 Ni 0.950 4253 0.404 3.800 2.460 0.457 <0.01 <0.05 0.061 Ni 0.969 328HB 51.5 0.3773.810 1.350 0.473 <0.01 <0.05 0.059 Ni 1.010 43 56 0.345 3.890 1.6400.470 0.012 <0.05 0.054 Ni 1.410 42 56 0.336 3.770 1.580 0.462 <0.01<0.05 0.055 Ni 1.580 42 55 0.409 3.750 1.360 0.451 <0.01 <0.05 0.060 Ni1.620 44 54.5 0.371 3.730 1.510 0.457 <0.01 <0.05 0.060 Ni 2.000 46 580.467 3.660 2.000 0.448 <0.01 <0.05 0.062 Ni 2.120 45 55 0.36 3.7-4 2.2<0.001 <0.02 <0.05 1.12 Ni 2.15   43.5 54 0.401 3.670 1.690 0.450 <0.01<0.05 0.062 Ni 2.560 395HB 53 0.367 3.660 1.460 0.463 <0.01 <0.05 0.060Ni 2.580 44 58 0.403 3.030 1.930 0.016 0.066 <0.05 0.145 Ni 2.840 44 560.336 3.040 1.930 0.012 0.061 0.103 0.149 Ni 2.870 40 51 0.240 2.9201.970 0.017 0.091 0.085 0.160 Ni 2.98 — — 0.383 3.35 1.92 <0.001 0.03270.119 0.117 Ni 2.98 42 53 0.350 3.020 2.070 0.018 0.094 0.080 0.150 Ni2.99 41 52 0.32 2.81 2.10 0.080 0.120 0.000 0.210 Cu, Ni 3.00   42.5 500.322 3.010 1.930 0.017 0.071 <0.05 0.144 Ni 3.010 38 50 0.32 3.13 1.90.030 0.07 0.13 0.17 Ni 3.04 39 50 0.340 3.100 1.990 0.016 0.120 <0.050.135 Ni 3.07 40 51 0.371 3.660 1.390 0.465 <0.01 <0.05 0.066 Ni 3.070409HB 55 0.402 3.060 2.100 0.020 0.085 <0.05 0.166 Ni 3.08 43 50 0.3843.080 2.130 0.016 0.074 0.088 0.158 Ni 3.08 338HB 49 0.32 2.92 1.750.030 0.1 0.14 0.16 Ni 3.1 40 49.5 0.384 3.090 2.080 0.019 0.079 0.1040.168 Ni 3.11 348HB 48 0.392 3.670 1.500 0.459 <0.01 <0.05 0.070 Ni3.190 44 58 0.240 3.20 2.39 0.050 0.070 0.010 0.240 Ni 3.21 38 49.50.392 3.63 2.52 0.0216 0.0832 0.0958 0.213 Ni 3.73   40.5 51 0.8 0.25<0.01 0 <0.01 1.59 1.98 —  40** 50 1.4 0.25 <0.01 3.0 <0.01 1.59 1.98 —   39.5** 49 0.8 0.25 <0.01 2.4 <0.01 1.59 1.98 —  42** 48.5 0.388 0.05<0.01 0.04 <0.01 1.5 1.56 Ni 0.06  320HB** 48 0.391 0.1 <0.01 0.03 0.041.62 1.61 Ni 1.15  43** 49 0.388 0.09 <0.01 0.05 2.08 1.43 1.53 Ni 0.07 42** 49 0.388 0.05 <0.01 0.02 0.01 1.52 1.61 Ni 0.05    40.5** 49 *Elements specified as other are present, otherwise indicated, in anamount of less than 2% **For these specific compositions, CVN was foundto be >40 J

1. A steel comprising a bainitic microstructure, comprising retainedaustenite, and/or carbide formers stronger than iron which are presentin the solid solution, wherein the microstructure comprises at least 50%vol. bainite.
 2. The steel according to claim 1, wherein the bainite isat least 50% high temperature bainite.
 3. The steel according to claim1, wherein the ferrite and/or perlite content is below 18%.
 4. The steelaccording to claim 1, wherein bainite is at least 76% vol.
 5. The steelaccording to claim 1, wherein the steel has a carbon content, wherein atleast 8% of the carbon content thereof is present in the form ofcarbides not belonging to the bainitic microstructure.
 6. The steelaccording to claim 1, wherein the bainite present comprises temperedbainite.
 7. The steel according to claim 1, characterized by a lowscattering structure characterized by a thermal diffusivity higher than8 mm2/s.
 8. The steel according to claim 1, comprising the followingcomposition, all percentages being indicated in weight percent: % Ceq =0.15-3.0 % C = 0.15-3.0 % N = 0-1.6 % B = 0-2.0 % Cr > 4.0 % Ni = 0-6.0% Si = 0-2.0 % Mn = 0-3 % Al = 0-2.5 % Mo = 0-15 % W = 0-15 % Ti = 0-2 %Ta = 0-3 % Zr = 0-3 % Hf = 0-3 % V = 0-12 % Nb = 0-3 % Cu = 0-2 % Co =0-6

The rest consisting on iron and trace elements, wherein% Ceq=% C+0.86*% N+10.2*% B
 9. The steel according to claim 8, wherein %V is above 0.1%.
 10. The steel according to claim 8, wherein 0% Cr ismore than 8%.
 11. The steel according to claim 1, comprising thefollowing composition, all percentages being indicated in weightpercent: % Ceq = 0.15-2.0 % C = 0.15-0.9 % N = 0-0.6 % B = 0-0.6 % Cr >11.0 % Ni = 0-12 % Si = 0-2.4 % Mn = 0-3 % Al = 0-2.5 % Mo = 0-10 % W =0-10 % Ti = 0-2 % Ta = 0-3 % Zr = 0-3 % Hf = 0-3 % V = 0-12 % Nb = 0-3 %Cu = 0-2 % Co = 0-12

The rest consisting on iron and trace elements, wherein% Ceq=% C+0.86*% N+1.2*% B
 12. The steel according to claim 11, wherein% Ceq is 0.22% or more.
 13. The steel according to claim 11 wherein % Cris above 12%.
 14. The steel according to claim 1, comprising thefollowing composition, all percentages being indicated in weightpercent: % Ceq = 0.2-0.9 % C = 0.2-0.9 % N = 0-0.6 % B = 0-0.6 % Cr =0.0-4.0 % Ni = 0-6.0 % Si = 0.2-2.8 % Mn = 0.2-3 % Al = 0-2.5 % Mo = 0-6% W = 0-8 % Ti = 0-2 % Ta = 0-2 % Zr = 0-2 % Hf = 0-2 % V = 0-4 % Nb =0-2 % Cu = 0-2 % Co = 0-6

The rest consisting on iron and trace elements, wherein% Ceq=% C+0.86*% N+1.2*% BCharacterized in that% Si+% Mn+% Ni+% Cr>2.0; or% Mo>1.2; or% B>2 ppm
 15. The steel according to claim 14, wherein % Si+% Mn+% Ni+%Mo is above 2%.
 16. The Steel according to claim 1, characterized inthat the sum of the amounts of those elements having an affinity forcarbon higher than iron selected from the group consisting of Cr, W, Mo,V, Ti, Nb, Ta, Zr and Hf is more than 4% in weight.
 17. The steelaccording to claim 1, characterized in that the microstructure comprisesless than 70% of alloyed carbides that can be attained with the chosencomposition.
 18. The Steel according to claim 1, characterized in thataccording to a tempering graph of the steel, the bainite present atempering degree which is smaller than that corresponding to a secondaryhardness peak, and a hardness of the steel is below a hardness level ofthe secondary hardness peak of the steel in an amount of at least 4 HRc.19. The Steel according to claim 1, characterized in that the bainitepresent comprises less than a 80% of a nominal % C of the steel.
 20. Thesteel according to claim 1, characterized in that the bainite presentcomprises less than a 80% of a nominal % C of the steel in an untemperedstate.